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Microscopic fracture mechanism of sintered high purity chromium
Materials Science and Engineering A285 (2000) 213 – 223 www.elsevier.com/locate/msea
Microscopic fracture mechanism of sintered high purity chromium Y. Matsumoto 1, S. Ohta, N. Aoki, M. Morinaga * Department of Materials Science and Engineering, Graduate School of Engineering, Nagoya Uni6ersity, Furo-cho, Chikusa-ku, Nagoya, Aichi 464 -8603, Japan
Abstract The microscopic fracture mechanism was examined experimentally for high purity chromium. With the specimens prepared by sintering and swaging, a series of tensile tests was carried out in various test conditions. The high purity chromium became more ductile as the test temperature increased, but the ductile-to-brittle transition temperature (DBTT) was not found clearly in the temperature range of 300–623 K. In addition, it was very brittle when tested in water, probably due to hydrogen embrittlement. From a fractographical analysis, it was shown that there were many micro-voids on the fracture surface, when the test temperature was less than 423 K. These micro-voids on a local region in the grain boundary seemed to grow up during deformation and to act as the initiation sites of intergranular fracture, and subsequently followed by transgranular cleavage fracture. The fracture strain increased with decreasing size of micro-voids, irrespective of the test conditions. The growth of the micro-voids was supposed to be one of the predominant factors for controlling the fracture of high purity sintered chromium. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Polycrystalline chromium; Microscopic fracture mechanism; Micro-voids; Environmental effect; DBTT
1. Introduction Polycrystalline chromium is brittle at room temperature. Its mechanical properties depend largely on alloying elements [1–5] and impurity elements [1 –3,6–11], and also on heat treatment processes [9 – 14]. For example, the presence of impurity elements such as nitrogen and oxygen makes chromium brittle [1 – 3,6 – 11]. The ductile-to-brittle transition temperature (DBTT) increases considerably as the content of such elements increases. Also, transgranular cleavage fracture initiates mainly on either grain boundaries or inclusions such as oxides and nitrides which precipitate on grain boundaries [6,7,13,14]. The mechanical properties of chromium are also influenced by surface states of specimens and test environments as well. According to our previous experiments, chromium is more brittle in air than in water [15–17]. In addition, chromium becomes more ductile * Corresponding author. E-mail address: [email protected] (M. Morinaga) 1 On leave from Department of Computer and Control Engineering, Oita National College of Technology, Maki, Oita 870-0152, Japan
in air when it is deformed in a cyclic way by repeating the loading and unloading process many times until it fails (hereafter, this deformation mode is called the multiple deformation mode) [17]. Thus, chromium possesses interesting mechanical properties, but most of them are not understood yet. In this paper, microscopic fracture mechanism of high purity chromium was examined in detail by tensile tests. With sintered polycrystalline chromium, tensile tests were performed in various conditions by adopting several test atmospheres, test temperatures, deformation modes and strain rates. The fractographical analyses were also made on the fracture surface using scanning electron microscopy.
2. Experimental procedure
2.1. Specimen preparation The chromium powders of higher purity than 99.93 mass% were sintered into a cylindrical shape in a hydrogen atmosphere at 1500 K for 7.2 ks under the applied stress of 118 MPa. The chemical composition is given in Table 1. The sintered chromium was then
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Table 1 Chemical composition of the specimens used in the present study Material
High purity chromium (powder) High purity chromium (as sintered)
Mass (%) C
H
O
N
S
P
Si
Al
Pb
Cu
Fe
Cr
0.026 0.004
0.004 0.0007
0.028 0.077
0.004 0.0037
0.0001 0.0018
0.0001 0.0034
0.006 0.002
0.001 0.002
0.0001 0.0001
0.0001 0.0022
0.003 0.024
99.93 99.88
swaged into a bar of 10 mm in diameter at the temperatures of 873–1173 K. The density of the specimen was higher than 96% of the bulk one. As shown in Fig. 1, the bar was machined into the shape of the tensile-test specimen. The gauge diameter was 2 mm and the gauge length was 3 mm. The specimen had a slender crystal grain, the size of which was about 1.5 mm in the longitudinal direction (i.e. the swaying direction), but about 100 mm in the transverse direction. The gauge part was mechanically polished with the emery papers and then with a buff, while dripping a solution of 0.3 mm A12O3 powders onto the specimen. Subsequently, in order to remove a work-hardened layer on the abrasive surface, electrolytic polishing was performed at room temperature by using a 5% HClO4 acetic acid solution.
2.2. Tensile test A tensile test machine used had an ultra-high vacuum chamber which was able to be evacuated as highly as about 10 − 7 Pa at room temperature. The cross-head speed of this tensile machine was variable in the range of 8.33×10 − 6 to 1.67 ×10 − 1 mm s − 1. The strain of
Fig. 1. Specimen for tensile test, and grain structures in the section: (a) normal to the direction of tensile axis and (b) parallel to the direction of tensile axis. Small black spots in these micrographs are etch pits.
specimens was measured at the gauge part using an optical device equipped with the image processing strain analyzer. The further detail of the tensile test machine is given elsewhere [18]. A series of tensile tests was performed in various conditions.
2.2.1. Test temperatures The test temperatures were varied in the range of 300–623 K. In particular, both 300 and 523 K were often employed in order to examine environmental effects and multiple deformation effects on the mechanical properties of chromium. 2.2.2. Strain rates The strain rate was set at 1.11 × 10 − 4 s − 1 in most cases. But the faster and the slower strain rates, 1.11× 10 − 2 and 2.78 × 10 − 6 s − l, were also employed in order to examine the strain rate dependence of the tensile properties. 2.2.3. Test en6ironments The test environments used were either air, or ultrahigh-vacuum (UHV) or water. For the test in air, a standard reference-atmosphere was prepared at first by passing the laboratory air through a distilled water vessel and then by charging it into the test chamber. As a result, this air contained a lot of moisture (4050%) at room temperature. For the test in UHV, vacuum in the chamber was controlled to be as low as 10 − 7 Pa at room temperature and 10 − 6 Pa at 623 K. For the test in water, the gauge part of the specimen was covered with a wet cloth. 2.2.4. Deformation modes In addition to the continuous deformation mode, the multiple deformation mode was employed for tensile tests. Here, the continuous deformation mode means a conventional deformation mode without any interruption of the test until the specimen fails. On the other hand, the multiple deformation mode means a cyclic deformation mode with repeated interruption during the test. Namely, the deformation at the first cycle was interrupted just after the specimen started yielding. Once unloaded, deformation was re-started, and interrupted again when the strain, xt, reached a set value in each cycle. Such stretching-unloading cycles were re-
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Fig. 2. Stress-strain curves for sintered chromium tested at 300 623 K in ultra-high vacuum.
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at 300 K and 33.8% at 623 K. In addition, necking of the specimen was seen in the test of higher than 473 K. The measured fracture strain changed gradually from 300 to 623 K as shown in Fig. 3, so that a ductile-tobrittle transition temperature (DBTT) was not clearly defined in this fracture strain versus temperature curve. In fact, there was even a small shoulder around 500 K, without showing any abrupt fracture strain change with temperatures. According to our previous results obtained from a small punch test [17], both the test environment and the multiple deformation mode were found to influence remarkably on the ductility of chromium when the test temperature was just below the DBTT. However, since the DBTT could not be clearly defined in the present tensile test, the test temperatures were varied in the wide range of 300–523 K in the following experiments.
3.2. Effects of test en6ironment, multiple deformation mode and strain rate on the fracture strain
Fig. 3. Temperature dependence of fracture strain measured in ultrahigh vacuum.
peated until the specimen fractured. Here, the total strain xt, was given by, xt =xc. · n, where xc is an increment of the strain in each cycle, and n is the number of cycles. The value of xc was empirically determined to be either 1.5 or 0.75% in the present experiment.
3. Results
3.1. Change in the fracture strain with temperature The measured stress – strain curves are shown in Fig. 2 for the sintered high purity chromium. These tests were performed at the five temperatures from 300 to 623 K in UHV of 10 − 7 10 − 6 Pa. Also, a continuous deformation mode was employed while keeping a constant strain rate of 1.11×10 − 4 s − 1. As usual, high purity chromium became more ductile with increasing temperature. For example, the fracture strain was about 1.5%
The stress–strain curves measured at 300 K are shown in Fig. 4. Here, the tensile tests were done either in air or in UHV employing both the continuous deformation mode and the multiple deformation mode. Also, in the latter deformation mode, the increment of the strain, xc, was set at either 1.5 or 0.75%. A test in water was performed only by the continuous deformation mode. The strain rate adopted in these tests was 1.11 × 10 − 4 s − 1. The measured fracture strain was 3.1% in air as shown in (a) 1.5% in UHV as shown in (d) and 0.2% in water as shown in (g). The result in water was very different from that obtained by the small punch test [17], since the fracture strain of the small punch test was larger in water than in air. The reason for this discrepancy is unknown at the moment. Besides, as shown in Fig. 4(d, e, f), upper and lower yield points were clearly observed in the stress-strain curve for the specimen tested in UHV, but not in air as shown in Fig. 4(a, b, c). Also, the yield stress was higher in UHV than in air. For the multiple deformation mode as shown in Fig. 4(b, c) or (e, f), the smaller strain increment, xc, caused the higher fracture strain, irrespective of the experiments in air and in UHV. In particular, when xc = 0.75%, the fracture strain was 3.1 times higher than the value obtained from the continuous deformation mode in air. In UHV it was 5.8 times higher because of the very small fracture strain of the specimen tested in the continuous deformation mode. On the other hand, such a multiple deformation effect on the fracture strain could not be examined in water because the specimen failed just after yielding, as shown in Fig. 4(g). The stress–strain curves measured at 523 K are shown in Fig. 5. In case of the continuous deformation
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Fig. 4. Stress-strain curves for sintered chromium at 300 K, tested (a – c) in air; (d – f) in ultra-high vacuum and (g) in water.
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Fig. 5. Stress-strain curves for sintered chromium at 523 K, tested (a, b) in air and (c, d) in ultra-high vacuum.
mode, the fracture strain in air was similar to that in UHV. However, yield points were clearly observed even at 523 K for the test in UHV, and the yield stress was higher in UHV than in air. In addition, there was no significant difference in the fracture strain between the continuous deformation mode and the multiple deformation mode at 523 K, in contrast to the result at 300 K shown in Fig. 4. Moreover, in order to examine the strain rate dependence of the fracture strain, the continuous deformation mode test was performed in air at 300 K with the three strain rates of 1.11×10 − 2, 1.11 × 10 − 4 and 2.78× 10 − 6 s − 1. The results are shown in Fig. 6. Plastic deformation scarcely occurred for the test at 1.11×10 − 2 s − 1. However, as the strain rate decreased from 1.11× 10 − 4 to 2.78×10 − 6 s − 1, the fracture strain increased from 3.1 to 8.0%. Also, the yield stress decreased gradually with decreasing strain rate; 350 MPa at 1.11 ×10 − 2 s − 1, 270 at 1.11 × 10 − 4 s − 1 and 220 at 2.78× 10 − 6 s − 1.
men tested at 300 K. The crack was extended radially from a fracture origin indicated by an arrow and a river pattern appeared on the surface. On the other hand, for the specimen tested at 623 K, as shown in Fig. 7(b), the trend of cleavage fracture still remained despite that the necking of the specimen yielded a large fracture strain of 33.2%. The region indicated by an arrow in Fig. 7(b) will be a sort of fracture origin. It is well known that fracture surface of polycrystalline chromium does not
3.3. Characteristics of fracture surface Typical fracture surfaces are shown in Fig. 7 for the two specimens tested in UHV at 300 and 623 K by the continuous deformation mode. As shown in Fig. 7(a), transgranular cleavage fracture took place in the speci-
Fig. 6. Strain rate dependence of stress – strain curves for sintered chromium at 300 K in air.
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Fig. 7. Cleavage fracture surfaces of sintered chromium tested at (a) 300 K and (b) 623 K in ultra-high vacuum at the strain rate of o; = 1.11 ×104 s − 1, and (c, d) microscopic voids in the intergranular rupture region shown in different magnifications. The fracture origin is indicated by an arrow in each figure, (a–c).
significant difference in the appearance of the fracture surface of all the specimens tested in the present experiment. However, one characteristic feature, depending strongly on the test conditions, was observed on a small area in the fracture origin. In Fig. 7(c, d), the small area in the vicinity of the fracture origin shown in Fig. 7(a) was enlarged and given again in larger magnifications. As shown in Fig. 7(d), a number of micro-voids were concentrated on the very small area of the fracture origin on the grain boundary. The occurrence of fracture associated with the growth and the linkage of micro-voids occurred in high purity sintered chromium. However, in this case, the micro-void did not grow up into a dimple, but remained in a round shape even after the specimen failed as shown in Fig. 7(d). It was expected that these voids themselves yielded the initiation sites of intergranular fracture. According to a series of previous investigations on the fracture of polycrystalline chromium [6,7,13,14], it is believed that transgranular cleavage fracture proceeds after intergranular fracture takes place locally in the grain boundary. This was also true in the present case, since the transgranular cleavage fracture initiated from the edge of micro-voids region, as shown in Fig. 7(d). However, in previous investigations, the fracture origin has been considered to be either oxides, nitrides or other inclusions precipitating on the grain boundaries [1–3,6– 11,13,14], whereas in the present study a small microvoid region of growing up on the grain boundary became a fracture origin. This discrepancy may be attributable partially to the purity of specimens. The purity of our specimen was very high, and hence the amount of precipitates such as oxides, nitrides and other inclusions was expected to be low in the grain boundary. Another reason may be due to the difference in the specimen preparation methods between them. Our specimen was prepared by sintering and swaging, so that the void formation was more or less promoted in it, compared to the specimen prepared by melting.
3.4. Nucleation and growth of 6oids
Fig. 8. Appearance of voids in sintered chromium: (a) small particles in a void; (b) slip lines in a void and (c, d) difference in the void growth between the neighboring grains.
exhibit any dimple structure at the DBTT or even higher temperatures [14,19], in agreement with the present observation. It was stressed here that there was no
It is well known that inclusions or second phase particles provide void nucleation sites when the applied tensile stress is large enough to decohere the particle/ matrix interface [20]. As explained earlier, the purity of chromium used in this study was extremely high, but still some particles with about 0.1 mm in diameter were observed, for example, they were seen inside a void as shown in Fig. 8(a). However, these small particles might not operate as nuclei to assist in forming micro-voids during the tensile test. Because, micro-voids of about 0.2 mm in diameter pre-existed in the specimen before the tensile test. Most of these observed voids may be generated in the course of sintering and swaging
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grown void. This implied that the void growth occurred in the slip process.
3.5. Test temperature dependence of fracture initiation sites
Fig. 9. Fracture origins of sintered chromium tested at (a) 300 and (b) 523 K. The intergranular fracture region is indicated by a dotted line in (a).
at high temperatures. The nucleation of voids without second phase particles during deformation has been reported even in single crystal chromium as well as polycrystalline chromium [21,22], but the appearance of the voids formed during deformation was very different from the present observation. These pre-existing micro-voids grew up during the tensile test. This is because, such a highly concentrated void region is deformable more easily than the surrounding crystal region with large shear modulus. The growth rate of micro-voids was affected strongly by the test conditions, and the void size was varied from 0.2 to 3.0 mm. Also, the micro-voids grew up in different ways, depending on crystal orientation. For instance, as shown in Fig. 8(c), the direction of the void growth was different between the upper and the lower crystal grains, and also as shown in Fig. 8(d), the size of the voids was also different between the right and the left crystal grains, where G.B. means the grain boundary. In addition, as shown in Fig. 8(b) by arrows, even clear slip lines were observed on the inner-wall of a largely
In case of the continuous tensile test performed in UHV, fracture initiation sites seemed to change with test temperatures. At the temperatures of 300–423 K, as shown in Fig. 9(a), intergranular fracture was observed on the grain boundary where a large number of micro-voids with about 1.5 mm in diameter were concentrated. The fracture surface was nearly perpendicular to the direction of the applied tensile stress. On the other hand, as shown in Fig. 9(b), any micro-voids were not seen in the fracture sites when the test was done at the temperature of 473–623 K, and transgranular cleavage fracture started from some local region in the grain boundary. In this case, the fracture surface was not perpendicular to the direction of the applied tensile stress. A wedge-shaped crack was supposed to first grow up in the grain boundaries by the slip mechanism and then to propagate toward the neighboring grains. However, even at the temperatures higher than 473 K, there were micro-voids on the cleavage surface, but they never acted as fracture initiation sites.
3.6. Test en6ironment dependence of micro-6oid growth Fig. 10 shows the appearance of the micro-voids observed on the fracture surface of the specimens tested by the continuous deformation mode in three environments at 300 K. Micro-voids grew up considerably during deformation. Here, the ratio of the void area to the total intergranular rupture area was defined in order to estimate quantitatively, the area fraction of the micro-voids. For example, the total intergranular rupture region is a central part surrounded by a dotted line
Fig. 10. Micro-voids on the intergranular rupture region of the specimens tested (a) in air; (b) in water and (c) in ultra-high vacuum.
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deformation mode in air shown in Fig. 10(a). This trend was observed not only in air, but also in UHV.
3.8. Strain rate dependence of micro-6oid growth The void area ratio changed with strain rates as shown in Fig. 12. In this case, every test was carried out employing the continuous deformation mode in air at 300 K. The fracture strain increased with decreasing strain rate, and also with decreasing void area ratio.
4. Discussion
4.1. Fracture strain and 6oid area fraction Fig. 11. Micro-voids on the intergranular rupture region of the specimens tested by the multiple deformation mode with strain increments of (a) xc =1.5% and (b) xc = 0.75%.
In Fig. 13, the measured ratios of the void area to the total intergranular rupture area are plotted against the fracture strain for all the specimens tested from 300 to 423 K in the present experiment. The results of strain rate and test temperature dependence are shown in Fig. 13(a), and the results of test environment and deformation mode dependence are shown in Fig. 13(b). As
Fig. 12. Micro-voids on the intergranular rupture region of the specimens tested at the strain rates of (a) o; = 1.11× 10 − 2 s − 1 and (b) o; = 2.78× 10 − 6 s − 1.
in Fig. 9(a), where a lot of micro-voids are embedded. This void area ratio was obtained readily from the SEM micrograph utilizing an image analyzing technique. As shown in Fig. 10, the void area ratio changed in the range of 0.45 – 0.60, depending on the test environments. It was evident that the fracture strain increased as this ratio decreased.
3.7. Multiple deformation mode effect on micro-6oid growth Fig. 11 shows the appearance of micro-voids in the specimen tested by the multiple deformation mode in air at 300 K (see Fig. 4(b, c)). The void area ratio was 0.34 for xc =1.5% and 0.11 for xc =0.75%. Either value was smaller as compared to 0.45 for the continuous
Fig. 13. A relationship between the void density and the fracture strain. Experimental data are shown in (a) for the strain rates and the test temperatures and in (b) for the test environments and the deformation modes.
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Fig. 14. Changes in the stress-strain curves with the strain rates for the specimens tested in water. Inserted photograph shows the fracture origin in sintered chromium.
described earlier, the fracture strain indeed increased with decreasing void area ratio in any test conditions. Thus, the area fraction of micro-voids in the total intergranular rupture region was a predominant factor for controlling the fracture of high purity sintered chromium regardless of test conditions. In other words the measured mechanical properties appeared to be treated consistently in view of the growth of the microvoids during deformation. The effective suppression of the micro-void growth will lead to the large fracture strain and hence to the high ductility of sintered chromium.
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cleavage fracture. Therefore, it may be said that the growth of micro-voids on the grain boundary is a rate controlling step in the whole fracture mechanism. However, as the test temperature rises, dislocation nucleation is more enhanced and also dislocation motion is thermally activated, so that the neighboring chromium crystals tend to become more deformable. By contraries the deformation of the void region will be suppressed to some extent, because it was surrounded by deformable regions. Above 473 K the void region did not play an important role in the deformation. In this case, a part of dislocations was supposed to accumulate on a certain grain boundary because of the orientation mismatch between adjacent grains for dislocation motion and to form a wedge-shaped crack, resulting in the brittle fracture as shown in Fig. 9(b). Also, this difference in the mechanism between low and high temperatures may give a reason why DBTT was not clearly observed in the fracture strain versus temperature curve shown in Fig. 3. If there is no mechanism change with temperatures, the fracture strain may be varied with temperature following a dotted curve shown in Fig. 3. However, when the temperature was lower than 473 K, the measured fracture strain curve lay below the dotted curve. This further embrittleness may be attributable to the presence of micro-voids in the specimen, since they provided fracture initiation sites as far as the temperature was in the range of 300–423 K. This additional void effect made the DBTT obscure. If there were no voids in chromium, the fracture strain at 300–423 K may follow the dotted curve, and the DBTT will lie between 523 and 623 K, in agreement with the previous experimental result [24].
4.2. Embrittlement caused by micro-6oids As explained before, pre-existing micro-voids may be concentrated on the narrow region in the grain boundary of sintered chromium. Upon loading at the lower temperatures than 473 K, this micro-void region will deform more readily than the neighboring chromium crystals with large shear moduli. As the presence of individual voids may act to concentrate strain in their vicinity [23], strain-induced void growth will occur preferably in such a region, as is analogous to the observation in porous materials (e.g. powder fabricated Ti and Ti – 6Al – 4V) [23]. The attendant increase in the void fraction may lower the flow stress of the narrow region [23], and hence void growth will further progress. As a result, slip deformation will be localized and enhanced into the very narrow region with the high void density, resulting in the initiation of intergranular fracture in it [23]. As shown in Fig. 8(b), the appearance of fine slip lines in the wall of a largely grown void may support this interpretation. According to previous studies on the fracture of chromium [13], the intergranular rupture will cause the transgranular
4.3. En6ironmental effect on fracture As shown in Fig. 4 chromium became more brittle in water than in air. This embrittlement of chromium in water observed in the present tensile test may be concerned with hydrogen embrittlement. In general, it is believed that it is difficult for hydrogen to diffuse into chromium because of the very low affinity of chromium for hydrogen. However, according to our previous study of small punch tests hydrogen indeed embrittled chromium in a dry hydrogen gas atmosphere [15]. In fact, the fracture strain increased with increasing strain rate. This is because, deformation at a slow strain rate gave enough time for hydrogen to diffuse into the specimen from the surface, so that hydrogen embrittlement was more enhanced as the strain rate decreased. So, further tensile test was continued in water at a slow strain rate of 2.78× 10 − 6 s − 1 in order to get strain rate dependence of the fracture strain. The result is shown in Fig. 14 together with the result of the strain rate, 1.11× 10 − 4 s − 1.
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When compared these results in water with the corresponding ones in air shown in Fig. 6, it was apparent that embrittlement occurred in water, even in the case of the strain rate of 2.78× 10 − 6 s − 1. Nevertheless, according to the results shown in Fig. 14, the fracture strain simply increased with decreasing strain rate. This correlation was inconsistent with the one expected from hydrogen embrittlement. However, in water, chromium will form a certain oxide film or hydroxide film instantaneously on the surface, even if fresh metal surface was exposed by deformation. Such an oxide or hydroxide film formed on the surface may suppress hydrogen absorption into chromium. For this reason, the total amount of hydrogen to be absorbed into chromium during deformation was supposed to be low, resulting in the less strain rate dependence of the fracture strain. In addition, the fracture origin of chromium in water often existed in the vicinity of the outer side of the specimen, as indicated by an arrow in Fig. 14. Thus, chromium will be embrittled to some extent by the invasion of hydrogen from the surface. Furthermore, it is likely that hydrogen goes preferentially to the void, so that high-pressure hydrogen gas is generated in the void, resulting in the acceleration of the void growth. In fact, as shown in Figs. 10 and 13(b), the void size and void area ratio were larger in water than in air. As shown in Fig. 4, chromium was more brittle in UHV than in air. This reason still remained unknown. However, recalling that the yielding was more remarkable and the yield stress was higher in UHV than in air, surface oxide softening [25,26] appears to be operating in some way in chromium because of the possibility of existing a certain thicker oxide film in air than in UHV.
4.4. Effect of multiple deformation on fracture The ductility of chromium was improved in air when the tensile test was performed by the multiple deformation mode. However, as shown in Figs. 4 and 5, the stress–strain curve changed continuously without showing any stress gap in each unloading and reloading step. This result implied that the multiple deformation neither induced multiplication of dislocation sources nor released dislocations adhered to obstacles. In other words, the multiple deformation probably did not increase the number of mobile dislocations in the crystal. Nevertheless, as shown in Fig. 13(b), the multiple deformation increased the fracture strain while showing a low void density. The reason why the void growth was retarded by the multiple deformation is still unknown. But, repeated cycles of loading and unloading in the multiple deformation process may lower the stress-concentrated area in the crystal. Because of the interruption of deformation, specific voids in some area may not keep growing preferentially until the size became large enough to act as a fracture origin. Instead, moder-
ately grown voids tend to be dispersed in the crystal because of the possible increase in void growth sites in the multiple deformation process. This may lead to the retardation of the fracture by the multiple deformation mode. If this mechanism is operating, it is natural that the multiple deformation effect was not observed at 523 K, since the void did not operate any more as a fracture origin at such a high temperature.
5. Conclusion In order to examine experimentally a microscopic fracture mechanism of sintered high purity chromium in detail, a series of tensile tests was carried out in various test conditions. Compared to the conventional continuous deformation, the multiple deformation improved significantly the ductility of chromium in air and in ultra-high vacuum. But, chromium was very brittle in water probably due to hydrogen embrittlement. At 300 423 K micro-voids grew up on the grain boundaries and acted as fracture origins. The growth rate of the voids on the fracture surface was influenced by the test conditions. However, irrespective of test conditions, the fracture strain increased as the void density decreased. The void growth or the void density was found to be associated with the fracture mechanism of sintered chromium.
Acknowledgements The authors would like to express their sincere thanks to Dr Y. Murata of Nagoya University for his helpful discussion and to Dr T. Sakaki of Tosoh Corporation Ltd. for his providing us chromium test samples. This research was supported by the Grant-in-Aid for Scientific Research from the Ministry of Education, Science, Sports and Culture of Japan.
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[19] F.P. Bullen, F. Henderson, H.L. Wain, Phil. Mag. 9 (1964) 803. [20] H. Mughrabi, Plastic deformation and fracture of materials, in: R.W. Cahn, et al. (Eds.), Materials Science and Technology, VCH, Weinheim, 1993. [21] A.V. Sameljuk, A.D. Vasilev, S.A. Firstov, Int. J. Refract. Metals & Hard Mat. 14 (1996) 249. [22] A.D. Vasilev, Scan. Electron Microsc. 111 (1986) 917. [23] R.J. Bourcier, D.A. Koss, R.E. Smelser, O. Richmond, Acta Metall. 34 (1986) 2443. [24] M. Ohmori, A. Kaya, Y. Harada, F. Yoshida, M. Itoh, J. Japan Inst. Met. 52 (1988) 223. [25] R.D. Noebe, R. Gibara, Phase boundary effects on deformation of bcc metals at low temperatures, in: K. Bramanian, et al. (Eds.), Structure and Deformation of Boundaries, TMS, AIME, 1986, pp. 89 – 108. [26] V.K. Sethi, R. Gibala, Acta Met. 25 (1977) 321.
Microscopic fracture mechanism of sintered high purity chromium Y. Matsumoto 1, S. Ohta, N. Aoki, M. Morinaga * Department of Materials Science and Engineering, Graduate School of Engineering, Nagoya Uni6ersity, Furo-cho, Chikusa-ku, Nagoya, Aichi 464 -8603, Japan
Abstract The microscopic fracture mechanism was examined experimentally for high purity chromium. With the specimens prepared by sintering and swaging, a series of tensile tests was carried out in various test conditions. The high purity chromium became more ductile as the test temperature increased, but the ductile-to-brittle transition temperature (DBTT) was not found clearly in the temperature range of 300–623 K. In addition, it was very brittle when tested in water, probably due to hydrogen embrittlement. From a fractographical analysis, it was shown that there were many micro-voids on the fracture surface, when the test temperature was less than 423 K. These micro-voids on a local region in the grain boundary seemed to grow up during deformation and to act as the initiation sites of intergranular fracture, and subsequently followed by transgranular cleavage fracture. The fracture strain increased with decreasing size of micro-voids, irrespective of the test conditions. The growth of the micro-voids was supposed to be one of the predominant factors for controlling the fracture of high purity sintered chromium. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Polycrystalline chromium; Microscopic fracture mechanism; Micro-voids; Environmental effect; DBTT
1. Introduction Polycrystalline chromium is brittle at room temperature. Its mechanical properties depend largely on alloying elements [1–5] and impurity elements [1 –3,6–11], and also on heat treatment processes [9 – 14]. For example, the presence of impurity elements such as nitrogen and oxygen makes chromium brittle [1 – 3,6 – 11]. The ductile-to-brittle transition temperature (DBTT) increases considerably as the content of such elements increases. Also, transgranular cleavage fracture initiates mainly on either grain boundaries or inclusions such as oxides and nitrides which precipitate on grain boundaries [6,7,13,14]. The mechanical properties of chromium are also influenced by surface states of specimens and test environments as well. According to our previous experiments, chromium is more brittle in air than in water [15–17]. In addition, chromium becomes more ductile * Corresponding author. E-mail address: [email protected] (M. Morinaga) 1 On leave from Department of Computer and Control Engineering, Oita National College of Technology, Maki, Oita 870-0152, Japan
in air when it is deformed in a cyclic way by repeating the loading and unloading process many times until it fails (hereafter, this deformation mode is called the multiple deformation mode) [17]. Thus, chromium possesses interesting mechanical properties, but most of them are not understood yet. In this paper, microscopic fracture mechanism of high purity chromium was examined in detail by tensile tests. With sintered polycrystalline chromium, tensile tests were performed in various conditions by adopting several test atmospheres, test temperatures, deformation modes and strain rates. The fractographical analyses were also made on the fracture surface using scanning electron microscopy.
2. Experimental procedure
2.1. Specimen preparation The chromium powders of higher purity than 99.93 mass% were sintered into a cylindrical shape in a hydrogen atmosphere at 1500 K for 7.2 ks under the applied stress of 118 MPa. The chemical composition is given in Table 1. The sintered chromium was then
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Table 1 Chemical composition of the specimens used in the present study Material
High purity chromium (powder) High purity chromium (as sintered)
Mass (%) C
H
O
N
S
P
Si
Al
Pb
Cu
Fe
Cr
0.026 0.004
0.004 0.0007
0.028 0.077
0.004 0.0037
0.0001 0.0018
0.0001 0.0034
0.006 0.002
0.001 0.002
0.0001 0.0001
0.0001 0.0022
0.003 0.024
99.93 99.88
swaged into a bar of 10 mm in diameter at the temperatures of 873–1173 K. The density of the specimen was higher than 96% of the bulk one. As shown in Fig. 1, the bar was machined into the shape of the tensile-test specimen. The gauge diameter was 2 mm and the gauge length was 3 mm. The specimen had a slender crystal grain, the size of which was about 1.5 mm in the longitudinal direction (i.e. the swaying direction), but about 100 mm in the transverse direction. The gauge part was mechanically polished with the emery papers and then with a buff, while dripping a solution of 0.3 mm A12O3 powders onto the specimen. Subsequently, in order to remove a work-hardened layer on the abrasive surface, electrolytic polishing was performed at room temperature by using a 5% HClO4 acetic acid solution.
2.2. Tensile test A tensile test machine used had an ultra-high vacuum chamber which was able to be evacuated as highly as about 10 − 7 Pa at room temperature. The cross-head speed of this tensile machine was variable in the range of 8.33×10 − 6 to 1.67 ×10 − 1 mm s − 1. The strain of
Fig. 1. Specimen for tensile test, and grain structures in the section: (a) normal to the direction of tensile axis and (b) parallel to the direction of tensile axis. Small black spots in these micrographs are etch pits.
specimens was measured at the gauge part using an optical device equipped with the image processing strain analyzer. The further detail of the tensile test machine is given elsewhere [18]. A series of tensile tests was performed in various conditions.
2.2.1. Test temperatures The test temperatures were varied in the range of 300–623 K. In particular, both 300 and 523 K were often employed in order to examine environmental effects and multiple deformation effects on the mechanical properties of chromium. 2.2.2. Strain rates The strain rate was set at 1.11 × 10 − 4 s − 1 in most cases. But the faster and the slower strain rates, 1.11× 10 − 2 and 2.78 × 10 − 6 s − l, were also employed in order to examine the strain rate dependence of the tensile properties. 2.2.3. Test en6ironments The test environments used were either air, or ultrahigh-vacuum (UHV) or water. For the test in air, a standard reference-atmosphere was prepared at first by passing the laboratory air through a distilled water vessel and then by charging it into the test chamber. As a result, this air contained a lot of moisture (4050%) at room temperature. For the test in UHV, vacuum in the chamber was controlled to be as low as 10 − 7 Pa at room temperature and 10 − 6 Pa at 623 K. For the test in water, the gauge part of the specimen was covered with a wet cloth. 2.2.4. Deformation modes In addition to the continuous deformation mode, the multiple deformation mode was employed for tensile tests. Here, the continuous deformation mode means a conventional deformation mode without any interruption of the test until the specimen fails. On the other hand, the multiple deformation mode means a cyclic deformation mode with repeated interruption during the test. Namely, the deformation at the first cycle was interrupted just after the specimen started yielding. Once unloaded, deformation was re-started, and interrupted again when the strain, xt, reached a set value in each cycle. Such stretching-unloading cycles were re-
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Fig. 2. Stress-strain curves for sintered chromium tested at 300 623 K in ultra-high vacuum.
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at 300 K and 33.8% at 623 K. In addition, necking of the specimen was seen in the test of higher than 473 K. The measured fracture strain changed gradually from 300 to 623 K as shown in Fig. 3, so that a ductile-tobrittle transition temperature (DBTT) was not clearly defined in this fracture strain versus temperature curve. In fact, there was even a small shoulder around 500 K, without showing any abrupt fracture strain change with temperatures. According to our previous results obtained from a small punch test [17], both the test environment and the multiple deformation mode were found to influence remarkably on the ductility of chromium when the test temperature was just below the DBTT. However, since the DBTT could not be clearly defined in the present tensile test, the test temperatures were varied in the wide range of 300–523 K in the following experiments.
3.2. Effects of test en6ironment, multiple deformation mode and strain rate on the fracture strain
Fig. 3. Temperature dependence of fracture strain measured in ultrahigh vacuum.
peated until the specimen fractured. Here, the total strain xt, was given by, xt =xc. · n, where xc is an increment of the strain in each cycle, and n is the number of cycles. The value of xc was empirically determined to be either 1.5 or 0.75% in the present experiment.
3. Results
3.1. Change in the fracture strain with temperature The measured stress – strain curves are shown in Fig. 2 for the sintered high purity chromium. These tests were performed at the five temperatures from 300 to 623 K in UHV of 10 − 7 10 − 6 Pa. Also, a continuous deformation mode was employed while keeping a constant strain rate of 1.11×10 − 4 s − 1. As usual, high purity chromium became more ductile with increasing temperature. For example, the fracture strain was about 1.5%
The stress–strain curves measured at 300 K are shown in Fig. 4. Here, the tensile tests were done either in air or in UHV employing both the continuous deformation mode and the multiple deformation mode. Also, in the latter deformation mode, the increment of the strain, xc, was set at either 1.5 or 0.75%. A test in water was performed only by the continuous deformation mode. The strain rate adopted in these tests was 1.11 × 10 − 4 s − 1. The measured fracture strain was 3.1% in air as shown in (a) 1.5% in UHV as shown in (d) and 0.2% in water as shown in (g). The result in water was very different from that obtained by the small punch test [17], since the fracture strain of the small punch test was larger in water than in air. The reason for this discrepancy is unknown at the moment. Besides, as shown in Fig. 4(d, e, f), upper and lower yield points were clearly observed in the stress-strain curve for the specimen tested in UHV, but not in air as shown in Fig. 4(a, b, c). Also, the yield stress was higher in UHV than in air. For the multiple deformation mode as shown in Fig. 4(b, c) or (e, f), the smaller strain increment, xc, caused the higher fracture strain, irrespective of the experiments in air and in UHV. In particular, when xc = 0.75%, the fracture strain was 3.1 times higher than the value obtained from the continuous deformation mode in air. In UHV it was 5.8 times higher because of the very small fracture strain of the specimen tested in the continuous deformation mode. On the other hand, such a multiple deformation effect on the fracture strain could not be examined in water because the specimen failed just after yielding, as shown in Fig. 4(g). The stress–strain curves measured at 523 K are shown in Fig. 5. In case of the continuous deformation
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Fig. 4. Stress-strain curves for sintered chromium at 300 K, tested (a – c) in air; (d – f) in ultra-high vacuum and (g) in water.
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Fig. 5. Stress-strain curves for sintered chromium at 523 K, tested (a, b) in air and (c, d) in ultra-high vacuum.
mode, the fracture strain in air was similar to that in UHV. However, yield points were clearly observed even at 523 K for the test in UHV, and the yield stress was higher in UHV than in air. In addition, there was no significant difference in the fracture strain between the continuous deformation mode and the multiple deformation mode at 523 K, in contrast to the result at 300 K shown in Fig. 4. Moreover, in order to examine the strain rate dependence of the fracture strain, the continuous deformation mode test was performed in air at 300 K with the three strain rates of 1.11×10 − 2, 1.11 × 10 − 4 and 2.78× 10 − 6 s − 1. The results are shown in Fig. 6. Plastic deformation scarcely occurred for the test at 1.11×10 − 2 s − 1. However, as the strain rate decreased from 1.11× 10 − 4 to 2.78×10 − 6 s − 1, the fracture strain increased from 3.1 to 8.0%. Also, the yield stress decreased gradually with decreasing strain rate; 350 MPa at 1.11 ×10 − 2 s − 1, 270 at 1.11 × 10 − 4 s − 1 and 220 at 2.78× 10 − 6 s − 1.
men tested at 300 K. The crack was extended radially from a fracture origin indicated by an arrow and a river pattern appeared on the surface. On the other hand, for the specimen tested at 623 K, as shown in Fig. 7(b), the trend of cleavage fracture still remained despite that the necking of the specimen yielded a large fracture strain of 33.2%. The region indicated by an arrow in Fig. 7(b) will be a sort of fracture origin. It is well known that fracture surface of polycrystalline chromium does not
3.3. Characteristics of fracture surface Typical fracture surfaces are shown in Fig. 7 for the two specimens tested in UHV at 300 and 623 K by the continuous deformation mode. As shown in Fig. 7(a), transgranular cleavage fracture took place in the speci-
Fig. 6. Strain rate dependence of stress – strain curves for sintered chromium at 300 K in air.
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Fig. 7. Cleavage fracture surfaces of sintered chromium tested at (a) 300 K and (b) 623 K in ultra-high vacuum at the strain rate of o; = 1.11 ×104 s − 1, and (c, d) microscopic voids in the intergranular rupture region shown in different magnifications. The fracture origin is indicated by an arrow in each figure, (a–c).
significant difference in the appearance of the fracture surface of all the specimens tested in the present experiment. However, one characteristic feature, depending strongly on the test conditions, was observed on a small area in the fracture origin. In Fig. 7(c, d), the small area in the vicinity of the fracture origin shown in Fig. 7(a) was enlarged and given again in larger magnifications. As shown in Fig. 7(d), a number of micro-voids were concentrated on the very small area of the fracture origin on the grain boundary. The occurrence of fracture associated with the growth and the linkage of micro-voids occurred in high purity sintered chromium. However, in this case, the micro-void did not grow up into a dimple, but remained in a round shape even after the specimen failed as shown in Fig. 7(d). It was expected that these voids themselves yielded the initiation sites of intergranular fracture. According to a series of previous investigations on the fracture of polycrystalline chromium [6,7,13,14], it is believed that transgranular cleavage fracture proceeds after intergranular fracture takes place locally in the grain boundary. This was also true in the present case, since the transgranular cleavage fracture initiated from the edge of micro-voids region, as shown in Fig. 7(d). However, in previous investigations, the fracture origin has been considered to be either oxides, nitrides or other inclusions precipitating on the grain boundaries [1–3,6– 11,13,14], whereas in the present study a small microvoid region of growing up on the grain boundary became a fracture origin. This discrepancy may be attributable partially to the purity of specimens. The purity of our specimen was very high, and hence the amount of precipitates such as oxides, nitrides and other inclusions was expected to be low in the grain boundary. Another reason may be due to the difference in the specimen preparation methods between them. Our specimen was prepared by sintering and swaging, so that the void formation was more or less promoted in it, compared to the specimen prepared by melting.
3.4. Nucleation and growth of 6oids
Fig. 8. Appearance of voids in sintered chromium: (a) small particles in a void; (b) slip lines in a void and (c, d) difference in the void growth between the neighboring grains.
exhibit any dimple structure at the DBTT or even higher temperatures [14,19], in agreement with the present observation. It was stressed here that there was no
It is well known that inclusions or second phase particles provide void nucleation sites when the applied tensile stress is large enough to decohere the particle/ matrix interface [20]. As explained earlier, the purity of chromium used in this study was extremely high, but still some particles with about 0.1 mm in diameter were observed, for example, they were seen inside a void as shown in Fig. 8(a). However, these small particles might not operate as nuclei to assist in forming micro-voids during the tensile test. Because, micro-voids of about 0.2 mm in diameter pre-existed in the specimen before the tensile test. Most of these observed voids may be generated in the course of sintering and swaging
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grown void. This implied that the void growth occurred in the slip process.
3.5. Test temperature dependence of fracture initiation sites
Fig. 9. Fracture origins of sintered chromium tested at (a) 300 and (b) 523 K. The intergranular fracture region is indicated by a dotted line in (a).
at high temperatures. The nucleation of voids without second phase particles during deformation has been reported even in single crystal chromium as well as polycrystalline chromium [21,22], but the appearance of the voids formed during deformation was very different from the present observation. These pre-existing micro-voids grew up during the tensile test. This is because, such a highly concentrated void region is deformable more easily than the surrounding crystal region with large shear modulus. The growth rate of micro-voids was affected strongly by the test conditions, and the void size was varied from 0.2 to 3.0 mm. Also, the micro-voids grew up in different ways, depending on crystal orientation. For instance, as shown in Fig. 8(c), the direction of the void growth was different between the upper and the lower crystal grains, and also as shown in Fig. 8(d), the size of the voids was also different between the right and the left crystal grains, where G.B. means the grain boundary. In addition, as shown in Fig. 8(b) by arrows, even clear slip lines were observed on the inner-wall of a largely
In case of the continuous tensile test performed in UHV, fracture initiation sites seemed to change with test temperatures. At the temperatures of 300–423 K, as shown in Fig. 9(a), intergranular fracture was observed on the grain boundary where a large number of micro-voids with about 1.5 mm in diameter were concentrated. The fracture surface was nearly perpendicular to the direction of the applied tensile stress. On the other hand, as shown in Fig. 9(b), any micro-voids were not seen in the fracture sites when the test was done at the temperature of 473–623 K, and transgranular cleavage fracture started from some local region in the grain boundary. In this case, the fracture surface was not perpendicular to the direction of the applied tensile stress. A wedge-shaped crack was supposed to first grow up in the grain boundaries by the slip mechanism and then to propagate toward the neighboring grains. However, even at the temperatures higher than 473 K, there were micro-voids on the cleavage surface, but they never acted as fracture initiation sites.
3.6. Test en6ironment dependence of micro-6oid growth Fig. 10 shows the appearance of the micro-voids observed on the fracture surface of the specimens tested by the continuous deformation mode in three environments at 300 K. Micro-voids grew up considerably during deformation. Here, the ratio of the void area to the total intergranular rupture area was defined in order to estimate quantitatively, the area fraction of the micro-voids. For example, the total intergranular rupture region is a central part surrounded by a dotted line
Fig. 10. Micro-voids on the intergranular rupture region of the specimens tested (a) in air; (b) in water and (c) in ultra-high vacuum.
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deformation mode in air shown in Fig. 10(a). This trend was observed not only in air, but also in UHV.
3.8. Strain rate dependence of micro-6oid growth The void area ratio changed with strain rates as shown in Fig. 12. In this case, every test was carried out employing the continuous deformation mode in air at 300 K. The fracture strain increased with decreasing strain rate, and also with decreasing void area ratio.
4. Discussion
4.1. Fracture strain and 6oid area fraction Fig. 11. Micro-voids on the intergranular rupture region of the specimens tested by the multiple deformation mode with strain increments of (a) xc =1.5% and (b) xc = 0.75%.
In Fig. 13, the measured ratios of the void area to the total intergranular rupture area are plotted against the fracture strain for all the specimens tested from 300 to 423 K in the present experiment. The results of strain rate and test temperature dependence are shown in Fig. 13(a), and the results of test environment and deformation mode dependence are shown in Fig. 13(b). As
Fig. 12. Micro-voids on the intergranular rupture region of the specimens tested at the strain rates of (a) o; = 1.11× 10 − 2 s − 1 and (b) o; = 2.78× 10 − 6 s − 1.
in Fig. 9(a), where a lot of micro-voids are embedded. This void area ratio was obtained readily from the SEM micrograph utilizing an image analyzing technique. As shown in Fig. 10, the void area ratio changed in the range of 0.45 – 0.60, depending on the test environments. It was evident that the fracture strain increased as this ratio decreased.
3.7. Multiple deformation mode effect on micro-6oid growth Fig. 11 shows the appearance of micro-voids in the specimen tested by the multiple deformation mode in air at 300 K (see Fig. 4(b, c)). The void area ratio was 0.34 for xc =1.5% and 0.11 for xc =0.75%. Either value was smaller as compared to 0.45 for the continuous
Fig. 13. A relationship between the void density and the fracture strain. Experimental data are shown in (a) for the strain rates and the test temperatures and in (b) for the test environments and the deformation modes.
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Fig. 14. Changes in the stress-strain curves with the strain rates for the specimens tested in water. Inserted photograph shows the fracture origin in sintered chromium.
described earlier, the fracture strain indeed increased with decreasing void area ratio in any test conditions. Thus, the area fraction of micro-voids in the total intergranular rupture region was a predominant factor for controlling the fracture of high purity sintered chromium regardless of test conditions. In other words the measured mechanical properties appeared to be treated consistently in view of the growth of the microvoids during deformation. The effective suppression of the micro-void growth will lead to the large fracture strain and hence to the high ductility of sintered chromium.
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cleavage fracture. Therefore, it may be said that the growth of micro-voids on the grain boundary is a rate controlling step in the whole fracture mechanism. However, as the test temperature rises, dislocation nucleation is more enhanced and also dislocation motion is thermally activated, so that the neighboring chromium crystals tend to become more deformable. By contraries the deformation of the void region will be suppressed to some extent, because it was surrounded by deformable regions. Above 473 K the void region did not play an important role in the deformation. In this case, a part of dislocations was supposed to accumulate on a certain grain boundary because of the orientation mismatch between adjacent grains for dislocation motion and to form a wedge-shaped crack, resulting in the brittle fracture as shown in Fig. 9(b). Also, this difference in the mechanism between low and high temperatures may give a reason why DBTT was not clearly observed in the fracture strain versus temperature curve shown in Fig. 3. If there is no mechanism change with temperatures, the fracture strain may be varied with temperature following a dotted curve shown in Fig. 3. However, when the temperature was lower than 473 K, the measured fracture strain curve lay below the dotted curve. This further embrittleness may be attributable to the presence of micro-voids in the specimen, since they provided fracture initiation sites as far as the temperature was in the range of 300–423 K. This additional void effect made the DBTT obscure. If there were no voids in chromium, the fracture strain at 300–423 K may follow the dotted curve, and the DBTT will lie between 523 and 623 K, in agreement with the previous experimental result [24].
4.2. Embrittlement caused by micro-6oids As explained before, pre-existing micro-voids may be concentrated on the narrow region in the grain boundary of sintered chromium. Upon loading at the lower temperatures than 473 K, this micro-void region will deform more readily than the neighboring chromium crystals with large shear moduli. As the presence of individual voids may act to concentrate strain in their vicinity [23], strain-induced void growth will occur preferably in such a region, as is analogous to the observation in porous materials (e.g. powder fabricated Ti and Ti – 6Al – 4V) [23]. The attendant increase in the void fraction may lower the flow stress of the narrow region [23], and hence void growth will further progress. As a result, slip deformation will be localized and enhanced into the very narrow region with the high void density, resulting in the initiation of intergranular fracture in it [23]. As shown in Fig. 8(b), the appearance of fine slip lines in the wall of a largely grown void may support this interpretation. According to previous studies on the fracture of chromium [13], the intergranular rupture will cause the transgranular
4.3. En6ironmental effect on fracture As shown in Fig. 4 chromium became more brittle in water than in air. This embrittlement of chromium in water observed in the present tensile test may be concerned with hydrogen embrittlement. In general, it is believed that it is difficult for hydrogen to diffuse into chromium because of the very low affinity of chromium for hydrogen. However, according to our previous study of small punch tests hydrogen indeed embrittled chromium in a dry hydrogen gas atmosphere [15]. In fact, the fracture strain increased with increasing strain rate. This is because, deformation at a slow strain rate gave enough time for hydrogen to diffuse into the specimen from the surface, so that hydrogen embrittlement was more enhanced as the strain rate decreased. So, further tensile test was continued in water at a slow strain rate of 2.78× 10 − 6 s − 1 in order to get strain rate dependence of the fracture strain. The result is shown in Fig. 14 together with the result of the strain rate, 1.11× 10 − 4 s − 1.
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When compared these results in water with the corresponding ones in air shown in Fig. 6, it was apparent that embrittlement occurred in water, even in the case of the strain rate of 2.78× 10 − 6 s − 1. Nevertheless, according to the results shown in Fig. 14, the fracture strain simply increased with decreasing strain rate. This correlation was inconsistent with the one expected from hydrogen embrittlement. However, in water, chromium will form a certain oxide film or hydroxide film instantaneously on the surface, even if fresh metal surface was exposed by deformation. Such an oxide or hydroxide film formed on the surface may suppress hydrogen absorption into chromium. For this reason, the total amount of hydrogen to be absorbed into chromium during deformation was supposed to be low, resulting in the less strain rate dependence of the fracture strain. In addition, the fracture origin of chromium in water often existed in the vicinity of the outer side of the specimen, as indicated by an arrow in Fig. 14. Thus, chromium will be embrittled to some extent by the invasion of hydrogen from the surface. Furthermore, it is likely that hydrogen goes preferentially to the void, so that high-pressure hydrogen gas is generated in the void, resulting in the acceleration of the void growth. In fact, as shown in Figs. 10 and 13(b), the void size and void area ratio were larger in water than in air. As shown in Fig. 4, chromium was more brittle in UHV than in air. This reason still remained unknown. However, recalling that the yielding was more remarkable and the yield stress was higher in UHV than in air, surface oxide softening [25,26] appears to be operating in some way in chromium because of the possibility of existing a certain thicker oxide film in air than in UHV.
4.4. Effect of multiple deformation on fracture The ductility of chromium was improved in air when the tensile test was performed by the multiple deformation mode. However, as shown in Figs. 4 and 5, the stress–strain curve changed continuously without showing any stress gap in each unloading and reloading step. This result implied that the multiple deformation neither induced multiplication of dislocation sources nor released dislocations adhered to obstacles. In other words, the multiple deformation probably did not increase the number of mobile dislocations in the crystal. Nevertheless, as shown in Fig. 13(b), the multiple deformation increased the fracture strain while showing a low void density. The reason why the void growth was retarded by the multiple deformation is still unknown. But, repeated cycles of loading and unloading in the multiple deformation process may lower the stress-concentrated area in the crystal. Because of the interruption of deformation, specific voids in some area may not keep growing preferentially until the size became large enough to act as a fracture origin. Instead, moder-
ately grown voids tend to be dispersed in the crystal because of the possible increase in void growth sites in the multiple deformation process. This may lead to the retardation of the fracture by the multiple deformation mode. If this mechanism is operating, it is natural that the multiple deformation effect was not observed at 523 K, since the void did not operate any more as a fracture origin at such a high temperature.
5. Conclusion In order to examine experimentally a microscopic fracture mechanism of sintered high purity chromium in detail, a series of tensile tests was carried out in various test conditions. Compared to the conventional continuous deformation, the multiple deformation improved significantly the ductility of chromium in air and in ultra-high vacuum. But, chromium was very brittle in water probably due to hydrogen embrittlement. At 300 423 K micro-voids grew up on the grain boundaries and acted as fracture origins. The growth rate of the voids on the fracture surface was influenced by the test conditions. However, irrespective of test conditions, the fracture strain increased as the void density decreased. The void growth or the void density was found to be associated with the fracture mechanism of sintered chromium.
Acknowledgements The authors would like to express their sincere thanks to Dr Y. Murata of Nagoya University for his helpful discussion and to Dr T. Sakaki of Tosoh Corporation Ltd. for his providing us chromium test samples. This research was supported by the Grant-in-Aid for Scientific Research from the Ministry of Education, Science, Sports and Culture of Japan.
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